Nanostructural Tailoring to Induce Flexibility in Thermoelectric

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Nanostructural tailoring to induce flexibility in thermoelectric CaCoO thin films Biplab Paul, Jun Lu, and Per Eklund ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b06301 • Publication Date (Web): 12 Jul 2017 Downloaded from http://pubs.acs.org on July 15, 2017

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ACS Applied Materials & Interfaces

Nanostructural tailoring to induce flexibility in thermoelectric Ca3Co4O9 thin films Biplab Paul,* Jun Lu, Per Eklund Thin Film Physics Division, Department of Physics, Chemistry, and Biology (IFM), Linköping University, SE-581 83 Linköping, Sweden Keywords: Ca3Co4O9, thermoelectrics, nanostructure, flexible film, transferable film Corresponding author. E-mail [email protected]

ABSTRACT Due to their inherent rigidity and brittleness, inorganic materials have seen limited use in flexible thermoelectric applications. On the other hand, for high output power density and stability, the use of inorganic materials is required. Here, we demonstrate a concept of fullyinorganic flexible thermoelectric thin films with Ca3Co4O9-on-mica. Ca3Co4O9 is promising not only due to its high Seebeck coefficient and good electrical conductivity but also important due to the abundance, low cost and nontoxicity of its constituent raw materials. We show a promising nanostructural-tailoring approach to induce flexibility in inorganic thin film materials, achieving flexibility in nanostructured Ca3Co4O9 thin films. The films were grown by thermally induced phase transformation from CaO-CoO thin films deposited by rfmagnetron reactive cosputtering from metallic targets of Ca and Co, to final phase of Ca3Co4O9 on mica substrate. The pattern of nanostructural evolution during solid state phase

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transformation is determined by surface energy and strain energy contributions, while different distributions of CaO and CoO phases in the as-deposited films promote different nanostructuring during phase transformation. Another interesting fact is that the Ca3Co4O9 film is transferable onto arbitrary flexible platform from parent mica substrate by etch free dry transfer. The highest thermoelectric power factor obtained is above 1 × 10-4 Wm-1K-2 in a wide temperature range, and thus showing low temperature applicability of this class of materials

1. INTRODUCTION Microscale electronic components tend to operate on battery power,1 which has limitations on lifetime and requirement for recharging. This is not desired for wearable devices, where a possible solution could be the scavenging of body heat for electrical power generation by flexible thermoelectric converters (TEC).2 However, for wearable and other flexible applications, a technology transformation is required from rigid thermoelectrics to flexible thermoelectrics.

Organic materials, due to their inherent flexibility, have been preferred over inorganic materials for this purpose. Extensive investigations have been done on organic materials,3–5 With high thermoelectric performance reported for conjugated polymer, PEDOT:PSS with a thermoelectric figure of merit, ZT, of 0.25.6 Despite the advantages from low material cost and solution-synthesis possibility, polymer materials typically have low output power density and stability.7,8 For high output power density and reliable performance over longer period of time, particularly in hostile environments, the use of inorganic materials is inevitable. However, it then becomes necessary to overcome the problem of material rigidity.


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Recently, there have been some investigations on developing flexible TEC based on inorganic materials.9,10 In those investigations the flexible platforms are used to hold the thermocouples of inorganic materials, and the legs of the thermocouples are subjected to temperature gradient in out-of-plane direction of the flexible platform. The disadvantage of such thermocouple arrangements (vertical arrangements), leg height of the thermocouples being in the micrometer range, the temperature gradient along the active materials is low, resulting in low output voltage from the modules. Further, maximum power output from a wearable thermoelectric device requires thermal matching between body skin and air, and for that 3 – 5 mm leg height is investigated to be appropriate.11 Achieving such leg height in a flexible module with vertical leg arrangements is quite challenging. An alternative option can be the lateral arrangement of thermocouples, where thickness of leg materials is not important as they are subjected to temperature gradient along their length, in parallel with substrate plane.12 The additional advantage of such arrangements is that large number of thermocouples can be accommodated in a small area. However, with such lateral arrangements of thermocouples both the substrate and the thin leg materials need to be mechanically flexible.

There are some attempts for developing flexible thermoelectric device with such lateral arrangements of thermocouples. For that, thin legs of inorganic materials are deposited on flexible polymer substrates by printing method, e.g. screen printing, inkjet printing, and dispenser printing.13–16 However, the problem with those printing techniques is that the low processing temperature of the film, restricted by low temperature sustainability of the polymer substrate, cause rough interfaces of the grains in the film, resulting in scattering of charge carriers and thus drastic reduction in electrical conductivity. To reduce the grain boundary scattering of charge carriers, the thermocouple legs can be deposited by sputter-deposition on flexible substrates.17, 18 However, the mechanical flexibility of the leg materials is still a



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challenge, which needs to be addressed by inducing mechanical flexibility in inorganic thin films but with no deterioration of their electronic properties. Recently, Zhou et al. developed carbon nanotube based flexible TEC for room temperature wearable applications,19 however its applicability above room temperature is not examined.

Tailoring structure at nanoscale can induce new mechanical properties in inorganic materials. For example, pristine Al2O3 is rigid in nature, but hierarchical nanoarchitecting has been reported to produce squeezable Al2O3, with 50 % recoverability.20 Nanostructural engineering has been used to tailor the electronic and phononic properties of inorganic thermoelectric materials for the enhancement of their thermoelectric efficiency.21–25 However, such experiments to induce mechanical flexibility in these materials are unexplored.

Here, we report the growth of flexible Ca3Co4O9 thin films on flexible mica substrate. A novel nanostructural tailoring approach is demonstrated to induce flexibility in Ca3Co4O9 thin films without significant effect on their electronic properties. Flexible Ca3Co4O9 films can be applicable in a wide temperature range from room temperature wearable applications to waste heat recovery from hot curved surfaces (e.g., hot pipes) and for applications in hostile environments. Thermoelectric performance of the investigated films have been evaluated in terms of their power factors. High power factor (=S2/ρ, where S is the Seebeck coefficient and

ρ is the electrical resistivity) is more important than low thermal conductivity to achieve a high output power,26 in particular, for low power applications, e.g. wearable applications. However, sustaining high power factor in flexible materials comparable to their pristine bulk values is quite challenging. The formation of nanolaminar platelets is typical of Ca3Co4O9, because of its inherently layered structure. We show that the size and orientation of those platelet like grains can be controlled to achieve flexible mechanical properties of the films


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without compromising with their thermoelectric performance. The nanostructured Ca3Co4O9 films are produced by thermally induced phase transformation from CaO-CoO thin films deposited on mica substrates by reactive rf-magnetron cosputtering from elemental targets of Ca and Co. Muscovite mica is chosen as substrate as it can act as flexible substrate and at the same time can sustain high processing temperature of 700 °C. Muscovite mica forms a layered structure, where aluminosilicate layers being loosely bound by boundary layer of potassium (K+) ions it is bendable and easily be cleaved along the boundary layer. Further, the film is easily transferable from mica by dry transfer, i.e. mica can also act as sacrificial layer for the transferable film.

2. EXPERIMENTAL SECTION Ca3Co4O9 thin films were prepared by a two-step sputtering/annealing process. In this process, first CaO–CoO films were reactively cosputtered from Ca and Co targets onto the muscovite mica (00l) substrates by rf-magnetron sputtering at 0.27 Pa (2 mTorr) in an oxygen – argon mixture with oxygen 1.5 %, second the as-deposited films were annealed at 700 °C in O2 gas flow to form the final phase of Ca3Co4O9. Four series of samples, namely (Ts: 20 °C), (Ts: 225 °C), (Ts: 375 °C), (Ts: 675 °C), were deposited with varying substrate temperature from room temperature to 675 °C, but with same oxygen percentage (1.5 %) in the gas mixture. The total gas pressure during the sputtered deposition is kept low, 2 mTorr, so as to avoid the scattering of ionized species. On the other hand, to ensure the deposition rate ∼ 10 nm/min the oxygen content in the gas mixture is maintained at minimum possible value of 1.5 % before it oxidizes the elemental targets of Ca and Co. Above 1.5 % oxygen the surface of the Ca and Co targets is found to be rapidly oxidized, resulting in drastic reduction in deposition rate. The crystal structure and morphology of the films were characterized by θ – 2θ X-ray diffraction (XRD) analyses using monochromatic Cu Kα radiation (λ = 1.5406 Å), 5 ACS Paragon Plus Environment


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transmission electron microscopy by using a FEI Tecnai G2 TF20 UT instrument with a field emission gun operated at 200 kV and with a point resolution of 1.9 Å, and scanning electron microscopy (SEM, LEO 1550 Gemini). The θ– 2θ XRD scans were performed with a Philips PW 1820 diffractometer. The samples for TEM were prepared by ion beam thinning method. For cross-sectional TEM two pieces of the sample glued together face to face and clamped with a Ti grid and then polished down to 50 µm thickness. Finally, polished sample ion milled in a Gatan Precision Ion Polishing System (PIPS) at Ar+ energy of 5 kV and a gun angle of 5°, with a final polishing step with 2 kV Ar+ energy. For the TEM analysis of the film from its top, it was cut into 3 mm diameter disks, and then thin the sample to 50 µm. Finally, the samples were ion milled from the substrate side until electron transparent. The composition of the films was determined by EDS, with an accuracy ±5%. The temperature dependent inplane electrical resistivity and Seebeck coefficient were simultaneously characterized using an ULVAC-RIKO ZEM3 system in a low-pressure helium atmosphere.

3. RSULTS AND DISCUSSION 3.1. Structure of the Films


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Figure 1. Optical image of (a) as-deposited CaO-CoO film, (b) annealed Ca3Co4O9 film, (c) XRD pattern of post-annealed film (Ts: 20 °C).

Figure 1a shows an optical image of an as-deposited film, which was deposited with no substrate heating, i.e. the substrate was kept at room temperature 20 °C during sputtering deposition, and hence the film is denoted (Ts: 20 °C). Similarly, a series of other films (Ts: 225 °C), (Ts: 375 °C), and (Ts: 675 °C) are named after their deposition temperatures 225 °C, 375 °C, and 675 °C, respectively. The as-deposited film is yellowish in color. This appearance is similar for the rest of the samples (not shown). As-deposited films consist of CaO-CoO phases, which is consistent with observations on sapphire substrate.26 Figure 1b shows the post-annealed film (Ts: 20 °C) after annealing. After annealing, all the samples turn dark as shown in Figure 1b. This change in color is attributed to the phase transformation from CaO-CoO phase to final phase of Ca3Co4O9. In our previous study, we demonstrated three stage phase transformation to occur during annealing leading to the formation of final phase of Ca3Co4O9.27 7 ACS Paragon Plus Environment


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Figure 1c shows the θ– 2θ XRD scan for the post-annealed film (Ts: 20 °C). Broadened peaks at around 2θ = 8.66°, 17.51°, 26.62°, and 35.85° occur from (00l) planes of muscovite mica. Diffraction peaks at 2θ = 16.42°, 24.73°, and 33.25° are observed, originating from the (002), (003), and (004) planes of Ca3Co4O9. The peak from (001) plane of Ca3Co4O9 is not visible here as it coincides with the broadened peak of mica at 2θ = 8.66°. Apart from (00l) planes of Ca3Co4O9 one low intense peak from (-201) plane is visible in Figure 1c, which indicates the film (Ts: 20 °C) is not singly oriented, it has grains with mixed orientation. The XRD peaks in the θ– 2θ XRD scan of the annealed films (Ts: 225 °C), (Ts: 375 °C), (Ts: 675 °C) are so weak that almost coincide with the background (see Figure S-1 of Supporting Information). This is because, the orientation of Ca3Co4O9 film might not satisfy Bragg’s condition in the out-ofplane direction, which is consistent with the previous observation for CaCo4O9 film grown on SrTiO3 (111).27 The orientation and the crystal structure of those films were investigated by TEM and SEM and are discussed later.

Figure 2. SEM images of (a) as-deposited film (Ts: 20 °C), (b) as-deposited film (Ts: 225 °C), (c) as-deposited film (Ts: 375 °C), (d) as-deposited film (Ts: 675 °C), (e) post-annealed film


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(Ts: 20 °C), (f) post-annealed film (Ts: 225 °C), (g) post-annealed film (Ts: 375 °C), (h) postannealed film (Ts: 675 °C).

Figure 2a-d shows the SEM images of as-deposited films (Ts: 20 °C), (Ts: 225 °C), (Ts: 375°C), (Ts: 675 °C), and Figure 2e-h shows the SEM images of post-annealed films. The morphology of the as-deposited films changes from sample to sample, which is attributed to the different deposition temperature of the films. The morphology of the annealed films also varies from sample to sample. The formation of platelets like grains in the post-annealed films is evident from the Figure 2. Due to the inherently layered structure, the formation of nanolaminated platelets is typical for Ca3Co4O9. Controlling the size and orientation of those nanolaminar platelets is not trivial in the films.29-32 Here, we have modified the orientation of the nanolaminated grains in the films independently of substrate by controlling the growth condition. Figure 2a shows the SEM image of annealed film (Ts: 20 °C), showing both inplane and out-of-plane orientation of nanolaminated grains of Ca3Co4O9. In the films (Ts: 225 °C), (Ts: 375 °C) and (Ts: 675 °Cs) the nanolaminated grains tend to align nearly vertically (as shown in Figure 2f-h), i.e. the c-axis of the grains is along the in-plane direction of the sample. The phase of those films have been confirmed by TEM analyses and is discussed later. The thickness of the nanolaminated grains in the film (Ts: 225 °C) is found not to be uniform, certain distribution in grain thickness is evident form SEM image. The thickness of the nanolaminated grains in sample (Ts: 375 °C) is found to be almost uniform (also evident from Figure 2g) and estimated to be around 50 nm. When the deposition temperature is raised to 675 °C a distribution in grain thickness is observed in the film (Ts: 675 °C). SEM images of larger area of all as-deposited and post-annealed films are provided in Figure S-2 and S-3 in the Supporting Information.



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Different arrangements of the nanolaminated grains in the annealed films is due to the different self-arrangements of the grains during nanostructural evolution during phase transformation, which is likely to be influenced by the initial arrangements of CaO and CoO phases in the as-deposited CaO-CoO films. The energetic constraints that guide the selfarrangements is anticipated to include the surface and interface energy minimization, as well as strain-energy minimization. The self-assembly growth of layered cobaltate in chemical solution deposition (CSD) technique was studied before.28 The oriented growth of the films was explained as due to the external stress due to solvent evaporation. In another study, Fu et al. reported the c-orientation of the Ca3Co4O9 film grown on polycrystalline Al2O3 substrate by CSD technique,33 They argued that the interactive force of (00l) plane of Ca3Co4O9 with Al2O3 (00l) plane stronger than other planes,34 and so Ca3Co4O9 (00l) plane tend to nucleate onto the Al2O3 (00l) plane serving as seeds for c-axis oriented growth, resulting in the c-axis self-assembled orientation. So, substrate is believed to have stronger impact on selection of the film orientation. However, in our study the various orientation of the grains of polycrystalline Ca3Co4O9 on the same substrate for different deposition conditions negates the argument on substrate influence on texture selection of the film. In our case, it is rather so that the distribution of crystallographic orientations of the grains in a polycrystalline film evolve during post-deposition annealing through a number of kinetic processes.35 The final texture of a film depends on which texture-selection mechanism and driving force dominates. In the present case, the different arrangements of CaO and CoO nanophases in as-deposited films drive the strain force in different direction leading to the different nanostructure of the postannealed films.

Figure 3a shows a typical cross sectional TEM image of an annealed film (Ts: 675 °C). The near vertical orientation of nanolaminated grains in Figure 3a is consistent with the


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observation from SEM image analyses. The compositional analyses by EDS confirms the Ca to Co ratio to be Ca:Co ~ 0.73, which corresponds to Ca to Co ratio of Ca:Co = 0.75 in Ca3Co4O9. The presence of an amorphous layer of thickness ~ 50 nm between the substrate and the film is evident from Figure 3a. The amorphous layer is formed due to the high temperature treatment during annealing. In amorphous layer the presence of Ca (21.3 at. %) along with the elements from mica substrate, O (55.7 at. %), Al (7.9 at. %), Si (9.9 at. %), K (1.0 at. %), and Fe (4.2 at. %), is confirmed by EDS analyses, however with no trace of Co. The proportion of O, Al, Si, K, and Fe in amorphous layer is found to be equivalent to that of mica substrate. This indicates that at 700 °C, the layered structure of mica near the interfacial region collapses, forming an amorphous layer through absorption of Ca. The formation of such interfacial layer was confirmed for all the annealed films (not shown). Figure 3b shows top view of TEM image of annealed film (Ts: 675 °C). The grains are found to form a closed network, which is consistent with the observation from SEM analyses. Such network formation is desirable for avoiding any disruption of transport of charge carriers during flexible applications. The presence of void spaces between the grains is visible in the Figure. 3b, which indicates that the film (Ts: 675 °C) is not 100 % dense. This is consistent with the SEM observation in Figure 2h. A high resolution TEM (HRTEM) image in the inset of the Figure 3b shows the lattice imaging of the nanolaminated grains. From lattice imaging the interlayer spacing (d-spacing) of layered cobaltate is confirmed to be around 10.7 Å, which matches with the d-spacing for Ca3Co4O9. From TEM image analyses it is clear that the nanolaminated grains are not perfectly vertically aligned (see Figure 3a), i.e. c-axis of the grains makes certain angle of inclination (5 – 25°) with the substrate plane. Due to such outof-plane alignment of the nanolaminated grains Bragg’s condition is not satisfied and so the XRD peaks were weak in θ– 2θ XRD scan.



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Figure 3. (a) Cross sectional TEM image of a typical post-annealed sample (TS: 675 °C). The amorphous layer is of Ca (21.3 at. %), O (55.7 at. %), Al (7.9 at. %), Si (9.9 at. %), K (1.0 at. %), and Fe (4.2 at. %). (b) TEM image of post-annealed sample (TS: 675 °C) taken from top of the film surface.

3.2. Flexibility and Transferability of the Film

Figure 4. (a) Image of thin flexible film (TS: 675 °C), (b) demonstration of thin film preparation from the post-annealed film (TS: 675 °C).

Figure 4a shows a typical flexible film prepared from the sample (Ts: 675 °C). Figure 4b shows the different steps leading to thin flexible film. In step 1, the film is attached to a glass 12 ACS Paragon Plus Environment

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slide by wax with upside down condition. In step 2, the thickness of the mica substrate is reduced to 100 µm by physical delamination process. To further reduce the substrate thickness, mica is delaminated by sticky tape, as shown in step 3. After repeating such delamination several times, the substrate thickness is reduced to 20 µm, and the film looks like as shown in step 4. In step 5, the glass slide is kept in acetone for 12 h so that the wax is completely dissolved and a thin flexible film can be isolated. The film is bendable to a bending radius of 14 mm without any deterioration of its physical properties.

Figure 5a shows a cross sectional SEM image of the flexible film (TS: 675 °C). The inset of Figure 5a shows the magnified image of a small cross-sectional portion of the film. The substrate thickness is around 20 µm. Figure 5b shows a magnified image of the interfacial region of the film. The average film thickness is 250 nm. Inset of Figure 5b shows the magnified image of a small portion of the film. The arrangement of nanolaminated grains of Ca3Co4O9 is clearly visible in the image. Such grain arrangements enable the film (TS: 675 °C) to withstand higher stress (tensile and compressive stress) developed due to bending. The epitaxially grown oriented Ca3Co4O9 films do not allow such bending without developing cracks. Therefore, we grow polycrystalline films with nanolaminated grains with their c-axis randomly oriented in the sample plane (i.e., standing basal planes). This arrangement of the grains results in network formation with gaps between the grains (as shown in Figure 3b), which allow relative motion and grain boundary/dislocation glide during bending and thus sustain bending stress. In a fully dense oriented film, this relative motion is not possible, and thus develops cracks to release the bending stress. Observation by optical microscope confirms the absence of cracks in the film even after the repeated bending of the film to the bending radius of 14 mm in both directions. Figure 5c shows an optical image of a small area (3.2 × 2.4 mm2) of the film before and after bending. The film surface before bending is 13 ACS Paragon Plus Environment


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seemingly flat; however, after bending some local curvatures in the film are developed from the compressive stress due to bending (evident from Figure 5c), but with no crack on the surface of the bended film. Thus it is confirmed that the film is able to sustain both tensile stress and compressive stress when it is subjected to bending of radius 14 mm, and thus has no effect on thermoelectric properties of the film (see Section 3.3 for more details on thermoelectric properties). Due to similar grain arrangements the films (TS: 225 °C), (TS: 375 °C) have been found to withstand bending stress, when they are subjected to bending.

Figure 5. (a) Cross sectional SEM image of the flexible film (TS: 675 °C), (b) magnified cross sectional SEM image of the flexible film (TS: 675 °C), (c) optical image of small area of the film (TS: 675 °C) before and after bending.

Another promising aspect of the present study is the transferability of the film to other flexible platforms. With the emergence of flexible thermoelectrics, transfer of the films from rigid substrate onto flexible platform is a major challenge. Strategies like surface-energy-assisted transfer,36 water-penetration-assisted mechanical transfer,37 film transfer by using ultrasonic water bath,38 carrier-polymer-assisted transfer,39 have been demonstrated to transfer monolayer or few atomic layer of metal sulfide onto flexible polymer platforms, however, these strategies have not yet been examined on thick films (say thickness of several hundred nanometer). Recently, Lu et al. demonstrated the possibility of transfer of thick films by etching of sacrificial water-soluble layers.40 14 ACS Paragon Plus Environment

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In the present study, we examined if 250 nm thick Ca3Co4O9 film can be mechanically transferred from parent mica substrate to another flexible platform. The transferability of the film was examined by transferring the film on a sticky tape. The different steps of film transfer are shown in Figure 6a. In step 1, one side of the film is marked by a sharp blade. In second step, a tape was stick to the film. After that the film is isolated from the mica substrate by stripping (step 3). Figure 6b shows the optical image of the back surface of the mechanically stripped film. Some leftover thin mica layers are still sticking to the back surface of the stripped film (Figure 6b). This, however, does not affect the functionality of the film as one exposed surface is sufficient for thermoelectric energy converters.

Figure 6. (a) Demonstration of transformation of the film onto a sticky tape, (b) optical image of the back surface of the stripped film, (c) optical image of the small portion of the back surface of the stripped film.

However, the functionality of the film can be effected by the micro-cracks. Figure 6c shows the optical image of a small area (3.2 × 2.4 mm2) of the back surface of the stripped film. The presence of micro-cracks on the exposed part of the film is evident from the Figure 6c. To avoid micro-cracking we instead remove the mica from the substrate side in a similar way as demonstrated before in Figure 4a, however, with one exception. The film, instead of sticking to a glass slide, is adhered to a flat and sticky surface of a sticky tape (as shown in Figure 7a). Figure 7b shows the back surface of the film after removal of mica. Partial presence of mica 15 ACS Paragon Plus Environment


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layers is still observed as before, but no micro-cracks are observed by optical microscopic analyses. Figure 7c is a typical optical image of a small area (~ 3.2 × 2.4 mm2) of the back surface of the film, which shows no cracks on the film.

Figure 7. (a) Optical image of the film (TS: 675 °C) from substrate side after the film is adhered to a sticky tape, (b) image of the film after removal of mica from substrate side, (c) optical image of the small portion of the film after the removal of mica, (d) image of the bended film after it is transferred to the sticky tape.

3.3. Thermoelectric Properties Figure 8a shows the temperature dependent electrical resistivity of all the annealed films. The electrical resistivity of all the films does not vary much as a function of temperature until 250 °C. Beyond 250 °C rapid increase in electrical resistivity is observed with temperature, which is attributed to the release of oxygen from the films.41-44 The rate of increase in electrical resistivity of all the films is not same. Above 250 °C the ρ vs T curve of the film (TS: 20 °C) is much steeper than other films. A nominal increase in electrical resistivity with temperature of the film (TS: 675 °C) as compared to the other films is likely due to the fact that it prevents 16 ACS Paragon Plus Environment

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release of oxygen to a greater extent at high temperatures. The lower tendency of oxygen release from the film (TS: 675 °C) is attributed to its larger grain size than that of the film (TS: 375 °C) (see Figure 3g and h for comparison). This is because, the release of oxygen is more probable from the region near the grains’ surface, and with the increase in grains’ size the surface to volume ratio decreases, which in turn reduces the effect of oxygen release. Note that the electrical resistivity measurement was performed in a low pressure of helium gas, increasing the tendency of oxygen release at high temperatures. The oxygen release will be very limited at atmospheric condition. However, the main focus of the present study is the low temperature applicability of Ca3Co4O9 films, particularly for wearable applications, and hence high temperature stability is of limited importance. Near room temperature the electrical resistivity of the films (Ts: 20 °C), (TS: 225 °C), (TS: 375 °C), and (Ts: 675 °C), is 29.73, 25.00, 20.30, and 16.46 mΩcm, respectively. The highest electrical resistivity of the film (Ts: 20 °C) is attributed to both out-of-plane and in-plane orientation of the grains in the film. Due to the inherently layered structure the physical properties of Ca3Co4O9 is anisotropic in nature. The electrical resistivity along the c-direction of Ca3Co4O9 is higher than in (a, b)-plane. Due to both out-of-plane and in-plane orientation of the grains, the resistivity of the film (Ts: 20 °C) is higher. Room temperature value of electrical resistivity of the film (TS: 675 °C) although several times higher than textured Ca3Co4O9 thin films,45 it is comparable with the values obtained from undoped polycrystalline bulk Ca3Co4O9.46-51 As previously mentioned, the vertical arrangement of the nanolaminated grains of Ca3Co4O9 is favorable for flexible applications. Repeated bending (100 times) of the film (TS: 675 °C) shows no deterioration of electrical conductivity. It was same for the films (TS: 225 °C) and (TS: 375 °C).



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Figure 8. Temperature dependent (a) electrical resistivity, (b) Seebeck coefficient, (c) power factor of all the films from room temperature to 400 °C before bending. Comparison graph of (d) electrical resistivity, (e) Seebeck coefficient, (f) power factor of the film (TS: 675 °C) before and after bending.

Figure 8b shows temperature dependent Seebeck coefficient (S) of all the annealed films. Seebeck coefficient of all the films varies with temperature following the same manner as electrical resistivity. The highest value of Seebeck coefficient near room temperature is obtained as 118 µV/K from the film (TS: 675 °C). Near room temperature, not much variation of Seebeck values is observed for different films, and remaining within 111 – 118 µV/K, which is comparable to the reported values for bulk Ca3Co4O9.47-49 Below 150 °C no considerable variation in Seebeck coefficient is observed among the films. Figure 8c shows the power factor (PF = S2/ρ) of all the films as a function of temperature. Near room temperature power factors ∼ 1 × 10-4 Wm-1K-2 are obtained from the films (TS: 675 °C), and achieving the highest value 1.18 × 10-4 Wm-1K-2 near 300 °C. The power factor in the film (TS: 675 °C), unlike bulk sample, is nearly flattened with temperature. It is remarkable for 18 ACS Paragon Plus Environment

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flexible oxide thin films to exhibit such high power factor near room temperature. Liu et al. demonstrated free-standing Ca3Co4O9/PEDOT-PSS composite thin film, however the room temperature value of power factor is almost seven times lower than the value obtained from flexible film (TS: 675 °C).52 Near room temperature the power factor of the flexible film (TS: 675 °C) is comparable to the values reported for undoped bulk polycrystalline Ca3Co4O9,49-52 and further enhancement of power factor of the film is possible by optimal doping.53, 54

To examine the bending effect on thermoelectric properties of flexible film (TS: 675 °C) its Seebeck measurement was performed after it was subjected to 100 times bending in both directions, however no notable change in the results is observed; whatever variation in Seebeck coefficient and electrical resistivity is found to be well below the error limit specified by ULVAC-RIKO ZEM3 system. Figure 8d-f compares the electrical resistivity, Seebeck coefficient, and power factor, respectively, of the film (TS: 675 °C) before and after bending. No remarkable change in the values of Seebeck coefficient, electricial resistivity, and power factor is observed. Small fluctuation in the values is well below the error limit. For flexible applications, flexibility of the substrate is necessary but not sufficient; the film need also to be flexible. Considering both the mechanical flexibility and thermoelectric properties, the presently developed Ca3Co4O9 films thus improve on other reports on flexible films.

4. CONCLUSION A fully-inorganic flexible film, Ca3Co4O9-on-mica, is developed. A nanostructural tailoring approach has been demonstrated to induce mechanical flexibility in Ca3Co4O9 thin films. Nanostructured Ca3Co4O9 film is obtained by thermally induced phase transformation from CaO-CoO thin film deposited on mica substrate by reactive rf-magnetron co-sputtering, to final phase of Ca3Co4O9. Mica acts as flexible substrate and at the same time as sacrificial 19 ACS Paragon Plus Environment


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layer for the film transfer onto other flexible platforms. The nanostructure of the film is influenced by the initial arrangements of CaO and CoO phases in the as-deposited films, which is controlled by controlling the deposition conditions: deposition temperature and percentage of oxygen in gas mixture. Flexible films are bendable to the bending radius of 14 mm without any deterioration of thermoelectric performance. The maximum power factor of the flexible film is 1.18 × 10-4 Wm-1K-2 near 300 °C, and does not change much as a function of temperature within the temperature range measured. With this high power factor and mechanical flexibility, the present films can be promising in the area of flexible thermoelectrics. Further enhancement of power factor is possible by optimal doping. The present approach can also be applicable to grow flexible films of other compounds in layered cobaltate family.

ASSOCIATED CONTENT Supporting Information θ– 2θ XRD scan for mica substrate and the post-annealed Ca3Co4O9 films (Ts: 20 °C), (Ts: 225 °C), (Ts: 375 °C), (Ts: 675 °C). SEM image of larger area of the films (Ts: 20 °C), and (Ts: 225 °C) before and after annealing. SEM image of larger area of the films (Ts: 375 °C), and (Ts: 675 °C) before and after annealing.

ACKNOWLEDGMENTS The research leading to these results has received funding from the European Research Council (ERC) under the European Community’s Seventh Framework Programme (FP/20072013)/ERC Grant Agreement No. 335383, the Swedish Research Council (VR) under Project No. 2012-4430, the Eurostars project E!8892 T-to-Power, and the Swedish Foundation for Strategic Research (SSF) through the Future Research Leaders 5 program. 20 ACS Paragon Plus Environment

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Nanostructural Tailoring to Induce Flexibility in Thermoelectric

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